The invention relates to liquid-phase-sintered SiC shaped bodies with improved fracture toughness and a high electrical resistance, and to methods for producing same.
Dense sintered SiC is distinguished by a combination of valuable properties, such as high hardness and wear resistance, ability to withstand high temperatures, high thermal conductivity, resistance to thermal shocks, and resistance to oxidation and corrosion. Due to these properties, the solid-state sintered SiC has nowadays become accepted as a virtually ideal material for sliding-contact bearings and axial face seals which are subject to wear in the chemical apparatus and mechanical engineering sectors. Alpha- or beta-SiC sintering powder with particle sizes in the submicron range (mean grain size  less than 1 xcexcm) and a simultaneous addition of up to 2% carbon and boron have been recognized as prerequisites for the pressureless sintering. As an alternative to boron or boron compounds, it is also possible to use aluminum and aluminum compounds or beryllium and beryllium compounds, the net result being the same. During the sintering, an SiC polytype transformation and grain growth take place, the extent of which processes are dependent on the nature and quantity of the sintering additives and the sintering temperature. Boron-doped beta-SiC powders tend toward secondary recrystallization (exaggerated grain growth), while starting from alpha-SiC powder with boron or aluminum doping results in a fine-grained globular or bimodal microstructure, comprising a fine matrix with plateletlike crystals showing a high aspect ratio. Generally, the solid-state sintered SIC bodies have a transcrystalline fracture mode and, at a relative density of 95-98% of the theoretical density (% TD), reach a room-temperature strength of up to 450 MPa, which is retained even at elevated temperatures of up to 1500xc2x0 C. However, the brittleness of this SiC ceramic represents an obstacle preventing any broader expansion of its use in industry. Owing to the high fraction of covalent bonding, SiC is extremely brittle, and even small flaws in the microstructure of the material may lead to sudden failure of a component.
Therefore, it is necessary to develop Sic ceramics with an improved fracture behavior and to meet the requirements imposed on the reliability of the components.
Nowadays, dense sintered Sic with an intercrystalline fracture mode, improved fracture toughness and room temperature strength can be produced in a similar manner to silicon nitride (Si3N4)using a liquid phase sintering process. The addition of suitable metal oxides or nitrides which form low-melting eutectics with SiC and the adhering SiO2 results, during the sintering operation, in the formation of a liquid phase which makes a decisive contribution to the densification of the ceramic. The recommended sintering additives are mainly mixtures of Al2O3 and Y2O3 (M. Omori et al.: U.S. Pat. No. 4,502,983 (1985) and U.S. Pat. No. 4,564,490 (1986), and R. A. Cutler et al.: U.S. Pat. No. 4,829,027 (1988)),but also Y2O3 or other sesquioxides of the rare earth metals (RE2O3) in combination with aluminum nitride (AlN) (K. Y. Chia etal.: U.S. Pat. No. 5,298,470 (1994), M. Nader, Dissertation: INAM, University of Stuttgart (1995), I. Wiedmann et al.:xe2x80x9cFlxc3xcssigphasensintern von Siliciumcarbidxe2x80x9d [Liquid-phase sintering of silicon carbide], in: Werkstoffwoche 1996, Symposium 7: Materialwiss. Grundlagen, Ed. F. Aldinger and H. Mughrabi, DGM Informationsgesellschaft, Oberursel (1997), 151-520, H. Kxc3x6lker et al.: DE 19730770 (1998)). The liquid-phase sintering may be carried out without pressure, i.e. under atmospheric pressure, if appropriate using a powder bed, or by employing the gas pressure sintering technique under elevated gas pressure without an embedding material. A liquid-phase 35 sintered SiC (LPSSiC) which has been introduced into the market by Wacker-Chemie GmbH under the brand name EKasic T exhibits a virtually pore-free, fine grained microstructure (mean SiC grain size approx. 1 xcexcm), a flexural strength of approx. 600 MPa, a fracture toughness which is 40 to 60% higher than that of solid state sintered SiC (SSiC) and a higher electrical insulating capacity (cf. Table 1).
The high electric resistance of EKasic T is caused by the presence of a continuous vitreous grain-boundary phase which surrounds the SIC grains in an insulating manner in the form of a thin film (approx. 1 nm). Together with the sintering aid Yxe2x80x94Al garnet, at the triple junctions of the SiC grains this continuous grain-boundary phase forms the so called binder phase, which joins the SiC hard-material grains to form a strong composite. Since the microstructural development and the properties of liquid phase-sintered SiC are decisively influenced by the selection of the composition (type and amount of the SiC and the binder phase) and the specific sintering parameters (gas atmosphere, pressure, temperature, time), it is not surprising that it is already possible for a number of SiC materials with very different properties to be produced from the sintering additive system AlNxe2x80x94Y2O3 (or RE2O3 or YAG).
According to the method which is known from DE 3344263 (corresponds to U.S. Pat. No. 4,569,922, Inv.: K. Suzuki/Asahi Glass), silicon carbide powders are sintered, together with sintering additives based on 3-30% by weight AlN, 0-15% by weight oxides of the IIIa transition metals (in particular yttrium, lanthanum and cerium) and 0-20% by weight SiO2, Al2O3 or Si3N4, under an argon or nitrogen atmosphere, without pressure or under gas pressure, at 2000-2200xc2x0 C. for from 2 to 10 hours to form SiC shaped bodies with a density of over 95% of the theoretically possible density, which exhibit flexural strengths of  greater than 800 MPa both at room temperature and at 14000xc2x0 C. The microstructure of these SiC sintered bodies exhibits elongate and/or platelike grains (mean grain length 3-5 xcexcm) of an SiCxe2x80x94AlN mixed crystal and a crystalline grain-boundary phase. The composition of the sintered bodies essentially comprises SiC with 2-20% by weight Al, 0.2-10% by weight N, 0.2-5% by weight 0 and 0 to 15% by weight of a metal from group IIIa. With sintering additives comprising less than 3% by weight AlN, only deficient sintered densities are obtained.
According to the process for the liquid-phase sintering of starting powder mixtures which, in addition to SiC, contain 1-10% by weight AlN, 1-15% by weight Y2O3 and up to 8% by weight SiO2 as sintering additives, which process is known from Japanese patent application No. 59-051384 (Publication No. 60-195057, 10.03.1985, Inventor: S. Nagano/Kyocera), sintered bodies with densities of between 95 and 99% TD which, due to their low resistivity of 0.5 ohmxc2x7cm, can be machined by spark erosion, are obtained after pressureless sintering for 2 3 hours under anargon atmosphere in the temperature range 1800-1950xc2x0 C.
According to the method which is known from U.S. Pat. No. 5,298,470,dated Mar. 29, 1994 (corresponds to EP 419,271, Inv.: Chia et al./Carborundum), once again starting powder mixtures which, in addition to silicon carbide, contain 0.5-15% by weight AlN, 0.1-15% by weight Y2O3 (or other RE2O3) and up to 10% by weight SiO2 as sintering additives are sintered without pressure under an argon atmosphere, preferably using a powder bed. The powder bed contains, in addition to SiC, the additives, which is intended to prevent the latter from becoming substantially depleted in the following sintered body, i.e. to prevent them from reacting with SiC to form gaseous decomposition products, which results in deficient densifying.
Claim 28 claims compact SiC sintered bodies with a high fracture toughness, with a microstructure characterized by, fine grains, homogeneity, a predominantly equiaxial grain shape and the presence of a discontinuous crystalline accompanying phase (preferably yttrium aluminate) in the interstices (triple junctions) of the Sic grains.
Claims 35 and 38 claim resistivities of less than 103 ohmxc2x7cm, predominantly clean SiC-SiC grain boundaries and fracture toughnesses of at least 7 MPaxc2x7m1/2, measured using the SENB method with notch widths of 0.5 mm.
Enriched levels of the elements aluminum and nitrogen can be detected in the Sic grains, but not in mixed crystal quantities(in the tenths of a percent to percent range), but rather insignificantly smaller quantities (200-300 ppm). According to the description of the process, this is not a case of pure liquid-phase sintering, but rather of mixed mode sintering, according to which in the initial phase liquid-phase sintering with particle rearrangement takes place, followed by solid state diffusion sintering.
As can be seen from column 17, lines 3-68 of U.S. Pat. No. 5,298,470, the liquid phasexe2x80x94apart from small discontinuous residual quantitiesxe2x80x94draws back to the triple junctions, and the second sintering stage, the solid state sintering, follows, making it possible to improve the properties of the sintered bodies.
According to column 11, lines 12 ff, the mixed mode sintering process may be carried out in a single stage, in the temperature range between 1900xc2x0 C. and 2050xc2x0 C. for three hours, or in two stages, for example one hour at 1900xc2x0 C. plus one hour at 2050xc2x0 C. For the pressureless mixed mode sintering process, the intention was obviously to produce fine-grained SiC sintered bodies with a high fracture toughness but without a continuous binder phase, since the prevailing opinion was evidently that the best properties of the end product can only be ensured with clean SiC-SiC grain boundaries.
However, with the proposed pressureless sintering methods, in particular in view of the long holding times and the comparatively high sintering temperatures, it is necessary to reckon with the known difficulties, such as a high loss of weight through the formation of volatile Al and Si suboxides (Al2O, AlO, SiO) and the formation of highly porous edge zones, which not only requires expensive further machining of the sintered bodies obtained, but also is linked to a high level of wear to the heating and insulating material of the furnace device and to scrap during sintering caused by faulty temperature control.
According to the process for producing electrically insulating SiC which is known from WO 97/03030 of 01.30.1997 (Inv.: D. Maravic/Negawatt), high resistivities of  greater than 107 ohmxc2x7cm are only achieved in SiC shaped bodies if more than 6 to 7% by volume of binder phase fractions remain in the finished sintered body following the liquid phase sintering of SiC with the addition of Y2O3+Al2O3. Owing to the high levels of sintering additives(10-12% by weight) required and the high losses of weight (2.8 to 4.5% by weight) which ensue during the pressureless sintering, this method is somewhat disadvantageous.
It has been reported by the Institut fxc3xcr Nichtmetallische Anorganische Materialien [Nonmetallic Inorganic Materials Institute] of the University of Stuttgart (Dissertation M. Nader 1995, I. Wiedmann et al.:xe2x80x9cFlxc3xcssigphasensintern von Siliciumkarbidxe2x80x9d [Liquid-phase sintering of silicon carbide], in: Werkstoffwoche 1996, Symposium 7: Materialwiss. Grundlagen, Ed. F. Aldinger and H. Mughrabi, DGM Informationsgesellschaft, Oberursel (1997), 151-520), that SiC with 10% by volume additives based on 60 mol % AlNxe2x80x9440 mol% Y2O3 can be completely densified without using a powder bed and without pressure at temperatures of 1950-2000xc2x0 C. under nitrogen or argon protective gas.
In view of the starting composition of 84.4 SiCxe2x80x9410.3 Y2O3xe2x80x943.0 AlNxe2x80x942.3 SiO2 (% by weight) and the loss of weight during sintering of  less than 3%, it is possible to conclude that there is at least 10% by weight of binder phase in the sintered bodies. The microstructure is fine-grained and presents globular SiC grains and a continuous binder phase which surrounds the SiC grains. Y2O3 and yttrium-aluminum-silicon oxynitride were detected in these bodies as crystalline binder-phase constituents, using X-ray diffraction analysis. These sintered bodies are in great need of improvement not only because of the large quantity of binder phase but also because of their macroscopic inhomogeneity (porous edge zones and segregation channels). Owing to very high mass losses, the sintering additives based on Al2O3xe2x80x94Y2O3 instead of AlNxe2x80x94Y2O3 require sintering to be carried out in the powder bed. According to the process which is known from DE 19730770 (1998, inv.: H. Kxc3x6lker et al./Elektroschmelzwerk Kempten GmbH), SiC starting powder mixtures containing 11.6 or 6.1% by weight sintering additives based on AlNxe2x80x94Y2O3 or AlNxe2x80x94YAG undergo gas pressure sintering under a nitrogen or argon atmosphere to form pore-free SiC bodies. As shown by transmission electron microscope examinations, it was possible to detect the binder phase in the sintered body as more than 10% amorphous or anamorphous film with a width of 15 xc3x85 in the Sic grain boundaries.
The fracture toughness, measured using the bridge method (sharp crack), at 3.0 MPaxc2x7m1/2, was only 20-40% higher than that of solid state sintered Sic (cf. Table 1, commercially available material variant EKasic D).
The production of even tougher SiC, i.e. the further improvement in the fracture toughness of liquid phase-sintered Sic, has recently been achieved by modification of the microstructure using so-called xe2x80x9cin situ platelet reinforcementxe2x80x9d. In this case, a two-stage sintering process or a high-temperature treatment of the sintered bodies following the sintering results in the formation of a microstructure with plateletlike sic grains. Due to the anisotropic grain growth of alpha-SiC at temperatures above the sintering temperature, the globular grains are redissolved to form plateletlike grains, the aspect ratio (ratio of platelet length to platelet thickness) can be influenced by varying the alpha-/beta-SiC content in the starting powder (cf. I. Wiedmann et al.: xe2x80x9cFlxc3xcssigphasensintern von Siliciumcarbidxe2x80x9d [Liquid-phase sintering of silicon carbide], in: Werkstoffwoche 1996, Symposium 7: Materialwiss. Grundlagen, Ed.F. Aldinger and H. Mughrabi, DGM Informationsgesellschaft,Oberursel (1997), 151-520).
However, platelet reinforcement is also possible starting from pure alpha-SiC powder (without any beta-SiC content). For example, if a commercially available EKasic T containing 5 to 6% by weight of binder phase and having a fracture toughness of 3.2 Mpaxc2x7m1/2 is annealed at 2050xc2x0 C. under an argon pressure of 20 bar for four hours, a 6Hxe2x86x924H SiC polytype transformation takes place during the recrystallization (cf. Table 2), and a homogeneous platelet microstructure with grain aspect ratios in the range from 2:1 to 10:1 is obtained. The mean platelet length, determined by image analysis, was 4.5 pm, the maximum platelet length approx. 20 pin. The K1c value determined using the bridge method was 4.0xc2x10.2 MPaxc2x7m1/2, i.e. the in situ platelet reinforcement led to a further increase in the fracture toughness of 25%.
According to the method which is known from U.S. Pat. No. 5,656,218 dated Aug. 12, 1997 (B. W. Lin et al./Industrial Technology Research Inst./Taiwan), the fracture toughness can also be increased by 22% by in situ platelet reinforcement using Sic sintering powder comprising alpha-SiC and beta-SiC in the weight ratio alpha-SiC/beta-SiC of 90:10 to 10:90, together with 5 to 20% by weight sintering additives based on Al2O3xe2x80x94Y2O3 and with two-stage sintering. In the two-stage sintering process, firstly consolidation by sintering is carried out at between 1800 and 1950xc2x0 C. for 0.5 to 8 hours, followed by annealing at 1900 to 2200xc2x0 C. for 0.1 to 4 hours. The fracture toughnesses, determined using the SENB method with an initial crack width of 0.15 mm, of a Sic specimen sintered with 3.8% by weight Y2O3 and 6.2% by weight Al2O3 were 4.9 MPaxc2x7m1/2 (fine grained globular microstructure) and 6.0 MPaxc2x7m1/2 (platelet microstructure with mean platelet lengths of  less than 5 xcexcm). Owing to the relatively high raw material and process costs, this method is as yet unsuitable for the mass production of liquid phase sintered Sic.
The good materials properties of the liquid phase sintered EKasic T have nowadays led to its use being preferred in applications which require hard, tough components which are able to withstand thermal shocks. Examples of such applications are dewatering elements in paper making machines, where the minimum possible wear has to be ensured in particular through a high edge stability of the ceramic plates (resistance to flaking). Furthermore, it is known that liquid-phase-sintered Sic is of particular importance, for example, as a functional ceramic for cooking hob hotplates or substrate plates for recording and reading heads and as a structural ceramic for high speed gas sealing rings or rolling contact bearings (balls, rolls, rings), since it has the following advantages over the liquid-phase-sintered silicon nitride which has hitherto been preferentially used in these application sectors:
SiC is more favorable in terms of powder costs, the thermal conductivity, the stiffness (modulus of elasticity) and the hardness of the SiC are higher, and the resistance to corrosion is significantly better.
However, the known liquid-phase-sintered SiC materials still present insufficient fracture toughnesses and an electric resistance which is too low compared to liquid phase sintered silicon nitride.
As is clear from the extensive prior art, it has not previously been possible to produce high density, liquid-phase-sintered SiC shaped bodies with a globular, fine grained microstructure and containing small amounts of binder phase at low cost while achieving a high fracture toughness of at least 4.0 MPaxc2x7m1/2 (measured using the bridge method) and high resistivities in the range from 107 to 1012 ohmxc2x7cm.
This results in the object of providing highly densified, virtually pore-free shaped bodies of liquid-phase sintered SiC which exhibit improved mechanical and electrical properties.
A further object of the invention is to specify a method with which shaped bodies of this nature having the desired properties can be produced economically and reproducibly in a simple manner, i.e. without using a two-stage high temperature sintering process for in situ microstructure reinforcement.
The first object is achieved by means of a 5 polycrystalline SiC shaped body which comprises 96% by weight to 99.5% by weight of a hard material phase, 0 to 0.1% by weight free carbon, and a partially crystalline binder phase, the hard-material phase comprising SiC and a Sixe2x80x94Cxe2x80x94Alxe2x80x94Oxe2x80x94N mixed crystal containing 0.2 to 1.5% by weight dissolved aluminum, 0.1 to 0.5% by weight dissolved nitrogen, 0.01 to 0.2% by weight dissolved oxygen, the hard-material phase being in the form of globular grains with a mean grain size of  less than 4 xcexcm, at least three grains enclosing a triple junction the globular grains having a structure with a core and a shell surrounding the core, the core comprising SiC and the shell comprising a Sixe2x80x94Cxe2x80x94Alxe2x80x94Oxe2x80x94N mixed crystal, the mixed crystal having an Al/N atomic ratio of 1.0:1.0 to 2.0:1.0, and the partially crystalline binder phase being formed from crystalline and amorphous phases of the rare earth xe2x80x94Alxe2x80x94Sixe2x80x94O quaternary system, the amorphous phase being present in the form of a grain boundary film which surrounds the globular grains, and the crystalline phase being present in the form of accumulations at the triple junctions containing rare earth aluminate.